Licentiate Thesis Licentiate Technology Production 18 No. 2017 Microstructure and mechanical and mechanical Microstructure cold of a 5 wt.% Cr properties of heat steel - Influence tool work treatment procedure. Muhammad Arbab Rehan
MUHAMMAD ARBAB REHAN MICROSTRUCTURE AND MECHANICAL PROPERTIES OF A 5 WT.% CR COLD WORK TOOL STEEL - INFLUENCE OF HEAT TREATMENT PROCEDURE. 2017 NO.18 ISBN 978-91-87531-56-9 (Printed) ISBN 978-91-87531-56-9 (Electronic) ISBN 978-91-87531-55-2
Tryck: Ineko AB, april 2017. Licentiate Thesis Production Technology 2017 No. 18
Microstructure and mechanical properties of a 5 wt.% Cr cold work tool steel - Influence of heat treatment procedure.
Muhammad Arbab Rehan
University West SE-46186 Trollhättan Sweden +46-520 22 30 00 www.hv.se
© Muhammad Arbab Rehan 2017 Print Book ISBN 978-91-87531-56-9 eBook ISBN 978-91-87531-55-2
Acknowledgements
I must admit that there are a lot of people who have helped me with the passage of time. However, I would like to acknowledge and express my gratitude to the most important ones, my parents. They have always been understanding, helpful and stood by me in all parts of life. Their positive approach and wise suggestions have always lifted me up. I would also like to thank my wife (Mehroze Nasim Khan) who has been supportive in all aspects of life. It has been a wonderful experience growing with her together with my beautiful daughter. Moreover, my sister and brothers who had given me a joyous childhood and have always contributed to my personality.
I would like to thank my colleagues in Material Science Department at Uddeholms AB for such a friendly environment especially during coffee breaks. Moreover, my industrial supervisors Berne Högman and Anna Medvedeva have always been supportive in listening to my ideas with patience and attention. Their advices are always valuable. Thank you for providing the support and help whenever I needed it. I would also like to thank Lars Göran Nord, Anders Thuvander, Per Erik Skogholm, Sebastian Ejnermark and Seshendra Karamchedu, who always have time for me for a discussion.
Finally, I would like to thank my academic supervisors Prof. Lars-Erik Svensson and Prof. Leif Karlsson. I have realized that it requires an immense patience to learn and teach someone. Without their guidance and support this work could not have been possible. I look forward to learn more from them.
I would also like to thank Uddeholms AB and K.K foundation for funding the research school of SiCoMap of which this industrial PhD is a part.
M. Arbab Rehan
11th of May 2017
iii
Populärvetenskaplig Sammanfattning
Titel: Mikrostrucktur och mekaniska egenskaper hos ett kallarbetsstål med 5% Cr- Inverkan av värmebehandling.
Nyckelord: Verktygstål; Austenitisering; Restaustenit; Värmebehandling; Mikrostruktur; Mekaniska egenskaper
Behovet av avancerade höghållfasta stål (AHSS) inom fordonsindustrin ökar dag för dag. Det motiveras främst av det faktum att AHSS kan användas som tunna plåtar och samtidigt ha hög hållfasthet. Detta gör det möjligt att reducera vikten hos bilar. Följaktligen ökar deras bränsleeffektivitet, vilket är positivt för miljön. Det förväntas också att AHSS inom en snar framtid kommer att ha ännu högre hållfasthet.
Verktygsstål med 5 vikts% Cr för kallbearbetning används ofta för att forma AHSS. Därför måste verktygsindustrin också följa med för att möta de utmaningar som AHSS ger i framtiden, det vill säga verktygsstålet måste ha en högre hårdhet, bättre tryckhållfasthet och ökad seghet. Ett sätt att uppnå målet är att utveckla en djupare förståelse för värmebehandlingen. Detta kan bidra till att förbättra de mekaniska egenskaperna hos verktygsstålet genom att modifiera processparametrarna vid värmebehandling av kallarbetsstål med 5 vikts% Cr.
Det konstaterades att en högre austenitiseringstemperatur kan användas för att åstadkomma en högre hårdhet, god tryckhållfasthet och tillräckligt seghet hos stålet. Emellertid kan en för hög austenitiseringstemperatur resultera i för kraftig förgrovning av de tidigare austenitkornen, vilket resulterar i låg slagseghet. Det konstaterades också att restaustenit kan omvandlas under härdning i två olika processer. För det första, anlöpning vid 525°C resulterade i omvandling av kvarvarande austenit till martensit vid kylning. För det andra, vid anlöpning vid 600°C under lång tid kan isotermisk omvandling av restaustenit till ferrit och karbider ske.
Dessa resultat användes för att förstå standardförfarandet för anlöpning av kallarbetsstål med 5 vikts.% Cr. Vidare diskuteras alternativa värmebehandlingsförfaranden med utgångspunkt från de viktigaste resultaten i denna avhandling.
v
Abstract
Title: Microstructure and mechanical properties of a 5 wt.% Cr cold work tool steel - Influence of heat treatment procedure. Keywords: Cold work tool steel; Heat treatment; Microstructural characterisation; Mechanical properties; Retained austenite. Print Book ISBN: 978-91-87531-56-9 eBook ISBN: 978-91-87531-55-2
The demand for Advanced High Strength Steel (AHSS) in the automotive industry is increasing day by day. It is mainly motivated by the fact that AHSS can be used as thin sheets while having high strengths. It enables weight reduction of the automobiles which consequently increases the fuel efficiency and has proven to be less harmful to the environment. It is also expected that AHSS will have even higher strength in the near future. Cold work tools steels with 5 wt.% Cr are commonly used to process AHSS. Therefore, the tool steel must meet the challenges in the future, i.e. have even higher hardness, compressive strength and toughness. One way of increasing the mechanical properties of the tool steel is by improving the heat treatment parameters. However, it is not possible without a deeper understanding of the heat treatment process. Therefore, this work presents investigations related to phase transformations occurring in a 5 wt.% Cr cold work tool steel during heat treatment. Furthermore, the influence of austenitisation and tempering temperatures on the microstructure and mechanical properties were investigated. The studies revealed that a higher austenitisation temperature can be used to achieve a higher hardness, good compressive strength and adequate toughness of the steel. However, too high austenitisation temperature may result in excessive coarsening of prior austenite grains which reduced the impact toughness. It was also found that retained austenite can transform during tempering by two different mechanisms. Firstly, when tempering at 525°C, carbides precipitate in retained austenite lowering its stability and permitting a transformation to martensite on cooling. Secondly, when tempering at 600°C for extended holding time retained austenite isothermally transforms to ferrite and carbides. This occurs by precipitation of carbides in retained austenite followed by a final transformation to ferrite and carbides. These results were used to understand the standard tempering procedure of the 5 wt.% Cr cold work tool steel. Furthermore, alternative heat treatment procedures are discussed based on the important findings presented in this thesis.
vii
Appended Publications
Paper A. Effect of austenitisation temperature and multiple tempering on the microstructure and impact toughness of a 5 wt.% Cr cold work tool steel.
M.A. Rehan, A. Medvedeva, L.E. Svensson and L. Karlsson, Proceedings of the 10th International Tooling Conference, Bratislava, Oct. 2016, 91-100.
I was the main author, designed the experiments and performed microscopy and thermodynamical calculations. Co-authors were involved in discussions of results and helped writing the manuscript.
Paper B. Effect of austenitisation and tempering on the microstructure and mechanical properties of a 5 wt.% Cr cold work tool steel.
M.A. Rehan, A. Medvedeva, B. Högman, L.E. Svensson and L. Karlsson, Steel Research Int., Vol. 87, Issue 12, Dec. 2016, 1609-1618.
Microscopy and thermodynamical calculations were performed by me. I was the main author of the manuscript. Co-authors contributed in designing the experiments, discussing the results and in writing of the manuscript.
Paper C. Retained austenite transformation during heat treatment of a 5 wt.% Cr cold work tool steel.
M.A. Rehan, A. Medvedeva, L.E. Svensson and L. Karlsson. Submitted to ‘Metallurgical and Materials Transactions A’.
I was the main author, designed the experiments, performed microscopy and thermodynamical calculations. Co-authors were involved in discussions of the results and in writing of the manuscript.
ix
Table of Contents
Acknowledgements ...... iii Populärvetenskaplig Sammanfattning ...... v Abstract ...... vii Appended Publications ...... ix Table of Contents ...... xi 1 Introduction ...... 15 1.1 Background ...... 16 1.2 Aim and approach ...... 17 1.3 Research questions ...... 18 2 Tool steels ...... 19 2.1 Cold work tool steels ...... 20 2.1.1 Trends in development of cold work tool steels ...... 21 2.1.2 Microstructure of 5 wt.% Cr cold work tool steels ...... 23 2.2 Production of tool steels ...... 23 2.2.1 Conventional production route ...... 23 2.2.1.1 Electric arc furnace ...... 23 2.2.1.2 Refining of the melt ...... 24 2.2.1.3 Up-hill casting ...... 25 2.2.2 Electro slag remelting ...... 26 2.3 Heat treatment of tool steels ...... 26 2.3.1 Soft annealing ...... 27 2.3.2 Hardening ...... 28 2.3.3 Tempering ...... 29 3 Experimental ...... 31 3.1 Material ...... 31 3.2 Heat treatments ...... 31 3.2.1 Single tempering treatments ...... 31 3.2.2 Double tempering treatments ...... 32
xi
3.2.3 Multiple tempering treatments ...... 32 3.3 Microscopy ...... 33 3.3.1 Light optical microscopy ...... 33 3.3.2 Scanning electron microscopy ...... 33 3.4 X-ray diffraction ...... 34 3.5 Dilatometry ...... 35 3.6 Thermodynamical calculations ...... 35 3.7 Mechanical testing ...... 36 3.7.1 Hardness ...... 36 3.7.2 Compressive strength ...... 36 3.7.3 Impact toughness ...... 36 4 Results ...... 39 4.1 Mechanical properties ...... 39 4.1.1 Hardness ...... 39 4.1.2 Compressive strength ...... 40 4.1.3 Impact toughness ...... 40 4.1.3.1 Fractography ...... 42 4.2 Microstructures ...... 43 4.2.1 Soft annealed microstructure ...... 43 4.2.2 As-quenched microstructures ...... 44 4.2.3 Tempered microstructures ...... 48 4.2.3.1 Single tempering treatments ...... 48 4.2.3.2 Double tempering at 200°C to 600°C ...... 50 4.2.3.3 Multiple tempering treatments at 525°C ...... 52 4.3 Transformation behaviour ...... 54 4.3.1 Hardening ...... 54 4.3.2 Single tempering ...... 56 4.3.3 Double tempering ...... 59 4.4 Thermodynamic calculations ...... 60
xii
5 Discussion ...... 63 5.1 Microstructure ...... 63 5.2 Effects of tempering ...... 64 5.2.1 Retained austenite transformation to martensite ...... 65 5.2.1.1 Single tempering treatments ...... 65 5.2.1.2 Multiple tempering treatments at 525°C ...... 67 5.2.2 Retained austenite transformation to ferrite and carbides 69 5.2.3 Mechanical properties ...... 70 5.2.3.1 Hardness and compressive strength ...... 70 5.2.3.2 Impact toughness ...... 71 5.3 Heat treatment recommendations ...... 73 5.4 Comments on research questions ...... 74 6 Conclusions ...... 77 7 Future work ...... 79 8 References ...... 81 9 Summaries of appended papers ...... 85 9.1 Paper A ...... 85 9.2 Paper B ...... 85 9.3 Paper C ...... 86 10 Appended Papers ...... 87 Paper A Effect of austenitisation temperature and multiple tempering on the microstructure and impact toughness of a 5 wt.% Cr cold work tool steel Paper B Effect of austenitisation and tempering on the microstructure and mechanical properties of a 5 wt.% Cr cold work tool steel Paper C Retained austenite transformation during heat treatment of a 5 wt.% Cr cold work tool steel
xiii
INTRODUCTION
1 Introduction
Manufacturing of steel is not a recent discovery. It has been traced back to around 1200 BC when martensitic steels were first used in weapons [1] and substituted the most commonly used metallic material ‘Bronze’. The introduction of carbon into the iron in combination with different heat treatments has been used throughout ages to provide steels with high hardness and wear resistance. This was all achieved without the help of any analytical instruments or the understanding of steel microstructure. Instead it was an art of craftsmanship. This craftsmanship remained unexplained until a pioneer German metallographer, ‘Adolf Martens’ described the phenomenon of the hardened microstructure [1].
Steel has played a vital role to enable economic growth and prosperity in society for example by supporting the creation of an infrastructure e.g. roads, railways, buildings, and bridges. Steel also holds an important position in fulfilling the future needs for development especially in the area of engineering. According to the World Steel Association, the monthly production of crude steel had increased to 132.4 million tonnes in November 2016 which was an increment of 5% from November 2015 [2]. This increase is mainly driven by fast developing countries like China which is producing approximately half of the world’s crude steel [2].
Steels have always been of interest for scientific research to understand the fundamental phenomena occurring in the microstructure. However, the goal of industrial research is rather to use the scientific knowledge to gain technical benefits. Optimising industrial processes for efficient production, fulfilling future demands for material development and understanding critical issues (e.g. failure mechanism) to achieve a longer life time of steel products are examples of such benefits. However, the task is not simple; the characterisation of microstructural constituents and their evolution during production is a primary challenge. A critical aspect is to understand the connection between the microstructure and the mechanical properties to provide possibilities to predict material behaviour.
Tool steels are fundamentally not different from other type of steels, i.e. they are alloys of iron and carbon. What makes them unique is that a wide range of alloying elements are used to provide the steel with characteristic properties for
15
INTRODUCTION optimum performance such as hardness, strength, wear resistance and toughness [1]. Tool steels are mostly used for industrial manufacturing of commercial products which makes them important in the field of production technology. There are many requirements on tool steels which demands care and optimization during manufacturing. However, the need for better and more advanced tool steels is fulfilled by exploring new alloy compositions as well as improved methods of production and processing.
1.1 Background
Cold work tool steels are used in industrial processes where the environment temperature is below 200°C. The common industrial applications of these steels are cutting, punching, pressing, drawing and rolling. There are many different variants of cold work tool steels as will be described in more detail in chapter 2. However, 5 wt.% Cr cold work tool steels are commonly used in processing Advanced High Strength Steels (AHSS) to manufacture body parts for automobiles like A-, B- and C- pillars etc. [3]. AHSS have yield strengths equal to or greater than 500 MPa and covers a variety of steels such as dual phase steels, martensitic steels etc. [4]. The extreme loading conditions on the tools when processing AHSS may lead to tool failures by plastic deformation, chipping, cracking or breakage of the tool [3].
The rapid increase in global population has caught the attention of the world. It is expected to have a number of negative effects such as depletion of natural resources, decreasing the quality of life and increasing emission of CO2 in the atmosphere. Therefore the idea of sustainability has been adopted in manufacturing industries and demands more environmental friendly products. In the automotive industry, there is an increasing demand for materials with low weight and high strength.
AHSS have high strength and good formability and are used as thin sheets (commonly 1-2 mm) to produce low weight parts for automobiles. This enables a reduction of the weight for lower fuel consumption and consequently lowering CO2 emissions [4, 5]. The high strength of these steels is a necessity not to compromise on safety during car accidents [5].
There has been an increasing use of AHSS in the automobile industry over the last few decades and the trend is expected to grow in the future. An example presented in fig.1, shows a forecasted increase in the use of AHSS for the North American light vehicle market. At the same time, there is a lot of ongoing research on AHSS aiming to achieve even higher strengths in the near future.
16
INTRODUCTION
Therefore, there are two main challenges for 5 wt. % Cr cold work tool steel regarding processing of AHSS.
x To fulfil future demands of achieving higher tool hardness and toughness. x To minimize the risk of tool failure to achieve a longer tool life.
180 160 140 120 100 80 60 40 20 Mass (kg) of AHSS per of per AHSS vehicle (kg) Mass 0 2012 2013 2014 2015 2016 2017 2018 2019 2020 2021 2022 2023 2024 2025
Figure 1 Forecast for utilisation of AHSS in North American light vehicle production [6].
1.2 Aim and approach
Heat treatment is the most critical process for tool steels as it can alter the properties of the tools significantly [7]. A non-optimised heat treatment can also contribute to early tool failure. Therefore, the aim of this work is to find ways of improving heat treatment practices for the 5 wt.% Cr cold work tool steel (Uddeholm Caldie) by investigating the relationship between austenitising temperature, tempering treatment, microstructure and mechanical properties.
It is known that 5 wt.% Cr cold work tool steels require a high hardness and good impact toughness, in particular for punching applications of AHSS. In addition, a high compressive strength is necessary to avoid plastic deformation. Wear resistance is another important mechanical property. However, it is believed that a higher wear resistance can be achieved by surface coating the tools. The wear resistance was therefore not investigated in this work.
17
INTRODUCTION
As a first step it was necessary to investigate the microstructure after heat treatment. Furthermore, to understand how this microstructure was formed, it was decided to analyse how the microstructure evolves during the heat treatment, i.e. what are the critical phase transformations involved. Finally, the resulting mechanical properties, i.e. hardness, impact toughness and compressive strength were evaluated.
1.3 Research questions
The final heat treatment consists of hardening (i.e austenitisation and quenching) and tempering. These processes affect the microstructure and thereby the mechanical properties of the steel. Thus the following research questions were considered the most important:
1. What are the effects of austenitisation temperature on microstructure and hardness? 2. How does the tempering temperature affect the microstructure and mechanical properties? 3. What are the effects of multiple tempering treatments on the microstructure and impact toughness? 4. How does holding time and temperature during tempering affect the transformation of retained austenite?
18
TOOL STEELS
2 Tool steels
Tool steels are the alloys used to manufacture the tools, dies and moulds that cut, shape and form other materials including steels, non-ferrous alloys, and plastics, at either elevated or room temperatures [1].
The main groups of tool steels are classified by American Iron and Steel Institute (AISI) as shown in Table 1.
Table 1 Classification of tool steels and their AISI designations [1].
Groups AISI Symbol Hot work tool steels Chromium or tungsten or molybdenum H High carbon and high chromium D Cold work tool steels Oil-hardened O Air hardening, medium alloy A High speed tool steels Tungsten T Molybdenum M Mold steels P Water hardening tool steels W Shock resistance tool steels S
It can be noticed that tool steels are arranged by their alloying composition (e.g. T and M), heat treatments (e.g. O, A and W), mechanical properties (e.g. S) or by application area (e.g H, P and D). In the early stage of development, tool steels were also identified by digit designations in the Unified Numbering System (UNS) for metals and alloys, established by Society of Automotive Engineers (SAE) and the American Society for Testing and Materials (ASTM). However, this system was merged into the AISI system where the last one or two digits of the AISI designation has its origin in the earlier UNS code. Examples of nominal chemical composition of some tool steels are provided in Table 2.
19
TOOL STEELS
Table 2 Nominal chemical composition of some commonly used tool steels (wt.%) [1, 8].
AISI C Mn Si Cr V W Mo designations H11 0.3 0.5 0.2 5.0 0.6 - 2.3 D2 1.5 0.4 0.3 11.3 0.8 - 0.8 A2 0.7 0.5 0.2 5.0 0.5 - 2.3 O1 0.9 1.0 0.3 0.5 0.3 0.5 - M2 0.8 0.4 0.3 4.0 2.0 6.0 5.0
2.1 Cold work tool steels
Cold work tool steels are used to make tools that are used in industrial applications, where the temperature is below 200°C. Cold work tools are used for example in automotive industry and paper industry. The characteristic mechanical properties for cold work tool steels are high hardness, high wear resistance, good toughness, and compressive strength [1, 3, 9, 10].
The main alloying elements of tool steels are carbon and carbide forming elements like Cr, Mo, V and W. The carbon content typically varies from 0.5 to 2.5 wt.% C and other (mainly carbide forming) alloying elements from 1 to 13 wt.%. Conventional cold work tool steels (A and D type) can be divided into three categories.
12 wt.% Cr cold work tool steels (1.5-2.2 wt.% C). 8 wt.% Cr cold work tool steels (0.5-1.5 wt.% C). 5 wt.% Cr cold work tool steels (0.5-1.5 wt.% C).
The chemical compositions of some cold work tool steels produced at Uddeholms AB using conventional methods are listed in Table 3. Also listed are two 5 wt.% Cr cold work tool steels (Caldie and Unimax) which are refined using the Electro Slag Remelting (ESR) process (described in section 2.2.2). There is also a variety of cold work tool steels produced by powder metallurgy (not discussed in this work).
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TOOL STEELS
Table 3 Conventional or ESR cold work tool steels produced by Uddeholms AB (wt.%) [3].
Uddeholm C Si Mn Cr Mo V W grades 12 wt.% Cr cold work tool steels Sverker 21 1.5 0.3 0.4 11.3 0.8 0.8 - Conventional Sverker 3 2.0 0.3 0.8 12.7 - - 1.1 Conventional 8 wt.% Cr cold work tool steels Sleipner 0.9 0.5 0.9 7.8 2.5 0.5 - Conventional Chipper 0.5 1.0 0.5 8.0 1.5 0.5 - Conventional 5 wt.% Cr cold work tool steels Calmax 0.6 0.3 0.8 4.5 0.5 0.2 - Conventional Caldie 0.7 0.2 0.5 5.0 2.3 0.5 - ESR Unimax 0.5 0.2 0.5 5.0 2.3 0.5 - ESR Rigor 1.0 0.3 0.6 5.3 1.1 0.2 - Conventional
2.1.1 Trends in development of cold work tool steels
High carbon (1.5-2.2 wt.%) 12 wt.% Cr cold work tool steels (e.g Sverker 21) are classical D type cold work tool steels. The high alloying content of these steels causes segregation of alloying elements during solidification, giving rise to significant formation of primary carbides in the steels. The typical primary carbides can be MC, M2C, M6C, M7C3 and M23C6. Primary carbides provide the steel with high hardness and wear resistance and remain largely undissolved during austenitisation. The hardened microstructure of Sverker 21 containing streaks of primary carbides is shown in fig. 2a. The martensitic microstructure provides the steel with high hardness while the primary carbides provide good wear resistance [9, 11].
The use of cold work tools is shifting towards automotive applications demanding tools with higher crack resistance. In 2001, a new cold work tool steel (Sleipner) was introduced with a lower amount of carbon and chromium than Sverker 21 (see Table 3). The intent was to lower the volume fraction of primary carbides to have a better resistance to cracking. It can be seen in fig. 2b that the hardened microstructure of Sleipner contains less streaks of primary carbides than Sverker 21. These steels have similar wear resistance as Sverker 21 but with better cracking resistance [12] and are an upgrade compared to Sverker 21.
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TOOL STEELS
a) Sverker 21
Streaks of primary carbides
10μm
b) Sleipner
Less streaks of primary carbides
10μm
c) Caldie
Undissolved carbides
10μm
Figure 2 Light optical micrographs showing the hardened microstructure of a) Sverker 21, b) Sleipner and c) Caldie. Notice that the streaks of primary carbides is lesser in Sleipner than in Sverker 21. The microstructure of Caldie contains few and small undissolved carbides. (Courtesy of Uddeholms AB).
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TOOL STEELS
In 2006, further steps were taken to produce cold work tool steels with even better cracking resistance. The idea was to have a cold work tool steel with a minimum amount of primary carbides. It was achieved by reducing the carbon content to 0.7 wt.% and chromium content to 5 wt.% (see Table 3). These steels have few and smaller primary carbides as shown in the hardened microstructure of Caldie in fig. 2c. The hardness of Caldie is similar to that of Sverker 21 but the cracking resistance is better than for Sleipner and Sverker 21.
It was shown that the lowering of alloying element content reduces the amount of primary carbides. Thus the trend in the development of cold work tool steels is aimed at having a steel microstructure where the presence of primary carbides does not introduce a risk for cracking.
2.1.2 Microstructure of 5 wt.% Cr cold work tool steels
The microstructure of 5 wt.% Cr cold work tool steels in the hardened condition contains some undissolved carbides in a mainly martensitic matrix (see fig. 2c). However some bainite and significant amounts of retained austenite can also be found depending on the selection of austenitisation temperature and cooling rate. The tempering treatment results in the transformation of retained austenite and precipitation of secondary carbides in martensite (discussed in section 2.3.3). In this way the steel gains ductility and obtains a good combination of hardness and dimension stability. Usual tempering temperatures for 5 wt. % cold work tool steel are 525°C or 540°C.
2.2 Production of tool steels
There are two common routes to produce tool steels i.e. conventional and powder metallurgy. Both of the routes consist of a series of different processes performed stepwise until a final product is ready.
2.2.1 Conventional production route
At Uddeholms AB, the conventional route is used to produce tool steels and is described in the following chapters.
2.2.1.1 Electric arc furnace
The first step in the production of tool steels is the melting of steel scrap in an electric arc furnace as shown in fig. 3. Graphite rods are used to form electric arcs that melt the scrap. Oxygen is blown into the furnace to promote oxidation
23
TOOL STEELS of impurities and oxidized elements float up to the top of the bath and are included in the slag. The chemical composition of the melt is balanced by additional alloying. The slag is removed when the molten steel is poured into a ladle.
a b
Figure 3 a) Steel scrap is melted in an electric arc furnace and the molten steel is poured into a ladle. b) Slag is removed from the top of the molten bath by re-pouring of the melt into a new ladle. (Courtesy of Uddeholms AB)
2.2.1.2 Refining of the melt
The melt is refined in a ladle furnace by slow heating of the melt with the help of graphite electrodes as shown in fig. 4a. Argon is blown into the melt for continuous stirring of the melt and refining takes place by deoxidation and desulfurisation. The chemical composition is checked and is further adjusted by the addition of alloying elements.
a b
Figure 4 a) Refining of the steel melt in a ladle furnace. b) Further refining of the steel by continuous stirring of the melt in a vacuum ladle degassing unit. (Courtesy of Uddeholms AB)
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TOOL STEELS
The ladle is then carried to a vacuum ladle degassing unit for further refining as shown in fig. 4b. In this process, argon gas is introduced at low pressures which enhances the stirring process. Continuous stirring at low pressure removes CO from the molten steel along with hydrogen, nitrogen, and sulphur.
2.2.1.3 Up-hill casting
At Uddeholms AB, the molten steel is carried to an up-hill casting unit for solidification as shown in fig. 5. The melt is poured into a tundish from where the melt is carried through narrow tubes to a casting plate and up into the mould where it solidifies.
Figure 5 The molten steel is solidified in an uphill casting unit. (Courtesy of Uddeholms AB)
During solidification, the wall of the mould is at lower temperature than the steel melt. Thus, the liquid in contact with the wall of the mould solidifies faster than the centre of the ingot giving rise to different types of grain structure of solidified steel. The last melt that solidifies have a different composition compared to the first solid fraction [13] which results in composition variations over the casted ingot. It is therefore a standard practice to perform a homogenization heat treatment on the ingot [14]. The standard homogenization treatment is carried out above 1200°C for 15 hours or more. In the case of ESR material the homogenization heat treatment is carried out after ESR.
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TOOL STEELS
2.2.2 Electro slag remelting
Steel grades that require extra cleanliness are processed with ESR as shown in fig. 6. In this process, the conventionally solidified ingot is used as an electrode and slag is placed at the bottom of the furnace. Heat is produced when a high AC current is passed from the electrode to the slag. Due to the high electrical resistivity of the slag, it melts first. The electrode starts melting when it is submerged in the molten bath of slag. The molten steel and the slag are contained in a copper mould which is cooled by water. The droplets of molten steel are denser than the slag and hence pass through it. They are collected in the pool of molten steel which solidifies with time. The highly reactive slag used in ESR removes the oxide inclusions and reduces the sulphur content [15]. The higher solidification rates in ESR than up-hill casting, reduces the carbide banding, carbide size and grain size [1, 15].
Figure 6 The ESR process remelts an uphill casted ingot with the help of a high AC current. (Courtesy of Uddeholms AB)
After ESR, the ingot is plastically deformed above the re-crystallisation temperature by forging or rolling. The primary reason is to break the coarse as- cast grains of the ingot in order to obtain a steel with a finer microstructure.
2.3 Heat treatment of tool steels
The most important heat treatments used in making of tools out of tool steels or providing the tools with desired properties are described in the following chapters.
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TOOL STEELS
2.3.1 Soft annealing
The goal of soft annealing heat treatment is to soften the material to enhance machinability and thus making it easier to shape the tool. For this reason, tool steels are delivered in soft annealed condition. From the microstructural point of view, soft annealing produces a microstructure with a ferritic matrix and spheroidised carbides. This not only softens the steel but also makes the microstructure uniform which is a precondition for performing a subsequent hardening treatment. Annealing temperatures can be just below, at or just above A1 temperature where the ferrite starts to transforms into austenite. The selection of temperatures for soft annealing depends on the desired properties for machining or on the composition of the steel.
The steel is held at the annealing temperature for some hours and slowly cooled to 500°C with cooling rates of typically 10 to 22 °C/h. The material is then air cooled to room temperature. A microstructure of Caldie soft annealed at 860°C for two hours is presented in fig. 7, showing spheroidised carbides in a ferritic matrix.
10 μm
Figure 7 A SEM micrograph of Caldie in soft annealed condition showing spheroidised carbides in a ferritic matrix. (Secondary electron image).
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TOOL STEELS
2.3.2 Hardening
Hardening is a process which changes the distribution of alloying elements and phases in a microstructure to make the steel harder. Hardening consists of three steps i.e. heating, austenitisation and quenching.
Austenitisation aims at reaching and holding at a temperature where austenite is the only stable phase. However, medium and high alloy tool steels contains carbides which are partly stable at higher temperatures. These steels are therefore austenitised at a temperature in the phase field of austenite and carbides. Care has to be taken with the selection of austenitisation temperatures, as inappropriate temperatures may cause coarsening of prior austenite grains which leads to poor toughness of the steel [16].
Quenching follows austenitisation and involves rapid cooling of the steel from the austenitisation to ambient temperature. The cooling rate is very critical as it determines the final phases in the as-quenched microstructure of the steel. The selection of cooling rates depend on the composition of the steel. For tool steels, rapid cooling (1 °C/s) will result in martensite formation while slower cooling may result in ferrite, pearlite and/or bainite formation.
During austenitisation primary carbides are partly dissolved enriching the austenitic matrix in alloying elements. A high alloy content of the austenite lowers the martensite start and finish temperature during quenching. For some steels the martensite start and finish temperature may be well above room temperature resulting in a fully martensitic microstructure. However, for medium and high alloyed steels, it is common that the martensite finish temperature is below the room temperature. Therefore, significant amounts of retained austenite can be found in the as-quenched microstructure.
As an example, a hardened microstructure of Caldie austenitised at 1050°C is shown in fig. 8. It shows prior austenite grains and undissolved carbides in the mainly martensitic matrix.
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TOOL STEELS
Prior austenite grains
Martensitic matrix
Undissolved carbides
10 μm
Figure 8 A SEM micrograph of Caldie after austenitisation at 1050°C. It shows prior austenite grains and undissolved carbides in mainly martensitic matrix. (Secondary electron image).
2.3.3 Tempering
Tempering is the final heat treatment after which the tool is ready to use. The mainly martensitic microstructure produced by hardening is hard and brittle. Thus the primary purpose of tempering is to increase the ductility of the steel. However, this must be done in a well-controlled manner i.e. without excessively decreasing the hardness. There is a wide range of temperatures (i.e. from 200°C to 600°C) that can be used for tempering. The selection of the tempering temperature is critical and depends on the desired properties or application area of the tool.
Tempering medium and high carbon steels at temperatures up to 200°C results in the precipitation of İ-carbides [16], while at temperatures between 200°C to 350°C cementite (M3C) is precipitated [7]. For tool steels, tempering between temperatures from 350°C to 600°C may result in coarsening of cementite and transformation of retained austenite. The temperature at which the retained austenite transforms depends on the steel composition.
Tempering in the range from 500°C to 600°C results in the precipitation of very small MC and M2C carbides (2-5 nm) [17, 18] in martensite. The carbides precipitate during holding at the tempering temperature. These precipitates are referred as ‘tempering carbides’. The formation of these nano sized particles in the microstructure increases hardness. This is known as ‘secondary hardening’ which produces the highest hardness at temperatures between 500°C to 525°C
29
TOOL STEELS for tool steels. However, tempering in the range from 540°C to 600°C results in coarser tempering carbides which contribute less to the hardness of the steel.
1050°C
0.5 h
1 °C/s 525°C 525°C Temperature (°C) Temperature 2 h Fan 2 h Fan cooling cooling
Time (h)
Figure 9 A schematic diagram showing the heat treatment of hardening and double tempering typically used for Uddeholm Caldie.
The transformation of retained austenite mainly occurs during cooling from the tempering temperature. For steels containing higher amounts of retained austenite (above 15 %), a single tempering treatment does not transform all the retained austenite. Therefore, double or sometimes triple tempering treatments are required to fully transform the retained austenite.
The tempering treatment is performed in cycles. Two or three cycles are the current standard recommendation by Uddeholms AB where each cycle consists of heating, holding (1 or 2 h) and cooling to ambient temperature. Fig. 9 shows an example of hardening at 1050°C and tempering twice at 525°C typically used for the steel Caldie.
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EXPERIMENTAL
3 Experimental
3.1 Material
An experimental batch of Uddeholm ESR Caldie was produced for this work. An Electro Slag Remelted ingot was hot rolled into a round billet with a diameter of 80 mm. The chemical composition (wt.%) of the steel is shown in Table 4.
Table 4 The chemical composition (wt.%) of the experimental batch of Uddeholm ESR Caldie.
C Si Mn P S Cr Mo V Al N 0.72 0.20 0.49 0.007 0.0002 5.04 2.33 0.5 0.009 0.006
3.2 Heat treatments
Three different sets of specimens were produced for the study. All specimens were taken from the centre of the soft annealed bar. They were first austenitised at temperatures of 1020°C, 1050°C or 1075°C (papers A and B). The specimens for paper C were austenitised only at 1050°C. Specimens were held at the austenitisation temperature for 30 minutes and then quenched to room temperature with a rate of 300 seconds between 800°C and 500°C.
3.2.1 Single tempering treatments
A first set of specimens (used in paper C) were heat treated in a dilatometer. The heat treatments were conducted to understand the retained austenite transformation during tempering by varying the holding time. All specimens after austenitisation were heated to 525°C or 600°C with a rate of 17°C/s. They were then isothermally tempered at 525°C or 600°C for different holding time followed by cooling with a linear rate of 1°C/s to room temperature (see Table 5). The name designated for all isothermally tempered specimens (T) shows a digit which represents the holding time (h).
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EXPERIMENTAL
Table 5 Dilatometry heat treated specimens (paper C).
Specimens Holding time (h)
As-quenched Q50 - T0.1 0.1 Austenitised at 1050°C T0.5 0.5
& T1 1
tempered at 525°C T2 2
for different holding times T6 6 T 10 10 Standard tempering T2 x 2 2 x 2 Tempered at 600°C T30 30 3.2.2 Double tempering treatments
Another set of specimens (used in paper B) were heat treated in a vacuum furnace. Tempering treatments were carried out on as-quenched specimens in a muffle furnace, twice for two hours at temperatures of 200°C, 250°C, 350°C, 450°C, 475°C, 500°C, 525°C, 540°C, 560°C or 600°C. The summary of the heat treatments are shown in Table 6.
Table 6 Specimens heat treated in a vacuum and muffle furnace (paper B).
Specimens Austenitisation Tempering temperature (°C) treatment 1020 - As-quenched 1050 - 1075 - Tempered (°C) 200, 250, 350, 450, 475, 1020, 1050, 1075 2 x 2 h 500, 525, 540, 560, 600.
3.2.3 Multiple tempering treatments
A third set of specimens (used in paper A) were produced to study the effect of austenitisation and number of tempering treatments at 525°C as listed in Table 7. The specimens for dilatometry were heat treated as described in section 3.2.1 while the specimen for microscopy and impact toughness testing were heat treated in vacuum and muffle furnaces (as explained in section 3.2.2).
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EXPERIMENTAL
These specimens were designated with Q for the as-quenched or T for the quenched and tempered conditions followed by two digits representing the austenitisation temperature (20 for 1020°C, 50 for 1050°C or 75 for 1075°C) and finally the number of tempering treatments at 525°C (for example x2 for a double tempering treatment).
Table 7 Specimens used to study the effects of austenitisation and multiple tempering at 525°C (paper A).
Sample Q20 T20x1 T20x2 Q50 T50x1 T50x2 T50x3 Q75 T75x1 T75x2 Aus. Temp. 1020°C 1050°C 1075°C Tempering treatment at 0 1 2 0 1 2 3 0 1 2 525 °C
3.3 Microscopy
Specimens for microscopy had dimensions of 17 x 17 x 10 mm (l x b x h, papers A and B). Microscopy was also performed on dilatometry specimens with dimensions of 10 x 4 mm (l x d, paper C). The specimens were ground with abrasive papers of mesh size 350, 500 and 1200. Later, the specimens were polished with diamond paste of 3 um and 1 um and etched with Picral (4% picric acid) followed by 2% Nital for revealing the microstructure. The etching times ranged from 5 s to 1 min depending on the heat treatment.
3.3.1 Light optical microscopy
A light Optical Microscope (LOM, Zeiss Axiophot) was used to measure the prior austenite grain size and for investigations of hardened and tempered microstructures. The average prior austenite grains size was measured at a magnification of 200x following the ASTM standard E-112 [19].
3.3.2 Scanning electron microscopy
A Scanning Electron Microscope (SEM, FEI Quanta 600 F) was used to study the microstructure at higher magnification. The SEM is equipped with a field emission gun and an Energy Dispersive X-ray Spectrometer (EDS, Oxfords Instruments). EDS was used to measure the chemical composition of the undissolved carbides. Secondary electron and back scattered electron detectors were used for microstructural investigations.
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EXPERIMENTAL
The electron beam in a SEM when interacts the specimen give rise to secondary electrons and backscattered electrons. The secondary electrons are emitted from a narrow depth (5-50 nm) and are used for higher resolution imaging. Backscattered electrons are emitted from a larger volume in the specimen. A material composed of elements with higher atomic numbers yield more backscattered electrons than a material containing elements with lower atomic number. This results in an atomic number contrast that can be used to differentiate between different phases such as different carbides in the steel. An acceleration voltage in the range between 10 kV to 20 kV was used.
3.4 X-ray diffraction
XRD was performed on microscopy and dilatometry specimens. The specimens were prepared in a similar fashion as specimens for microscopy except for the final etching, the specimens were electro-polished with a voltage of 40 V for 10 seconds in a Struers LectroPol-5 equipment.
XRD (Seifert 3003) was conducted to measure the volume fraction and the lattice parameters of retained austenite. XRD was carried out using unfiltered Cr Kα radiation. An acceleration voltage of 40 kV and a step size of 0.1° over the range of 2θ = 50°-165° were used. The software, Rayflex 2.407 (GE & Inspection Technologies) was used to measure the relative intensities of the diffracted peaks of martensite (110)α, (200)α, (211)α and the diffracted peaks of retained austenite (111)γ, (200)γ and (220)γ. The use of these peaks allows avoiding possible bias due to crystallographic texture [20]. The volume fraction of retained austenite was measured by performing curve fitting of the diffracted peaks on the plot of intensity as a function of 2θ. The amount of retained austenite was calculated by comparing the integrated intensities of the diffracted peaks of martensite and austenite according to Cullity [21]. The uncertainty in measuring retained austenite was 2%.
The lattice parameter (aγ) was calculated for individual peaks of austenite by plotting the lattice parameter as a function of cos2θ/sin2θ. The precise value of lattice parameter was obtained by extrapolating the linear fit against θ to 90°. The weighing of the data set is described elsewhere [22] where the highest weight was given to the largest value of 2θ. The standard deviation for the lattice parameter was calculated from the results of three measurements.
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EXPERIMENTAL
3.5 Dilatometry
A dilatometer is an instrument that is used to measure the volumetric changes in a specimen during a heat treatment. The volume changes are caused by thermal expansion and/or contractions and phase transformations in the material.
The specimens had a cylindrical shape with a length of 10 mm and a diameter of 4 mm. A push rod dilatometer (Dil 805 A/D, Bahr-Thermoanalyse) was used to study the phase transformations in the tool steel specimens. The dilatometer had a resolution of 0.05 µm / 0.05 °C (∆L/∆T). The specimens were heated with a high frequency inductive coil. Specimen temperatures were measured by a thermocouple, spot welded 1 mm apart on the middle of longitudinal surface of the specimen. The specimens had a cylindrical shape with a length of 10 mm and a diameter of 4 mm. Dilatometry data during quenching, isothermal tempering and cooling from the tempering temperature was used to construct three different types of plots:
i) Plots of length change against specimen temperature during cooling where transformation temperatures were evaluated by calculating the derivative of length change as a function of specimen temperature. The transformation start temperatures were defined as the initial deflection in the derivative. The standard error in measuring start and finish temperature was estimated to ± 5°C. ii) Plots comparing relative length change of specimens as a function of time during the first and second isothermal tempering treatments at 525°C. The relative length changes for different austenitisation temperatures were plotted together in a graph with a numberless y-axis for ease of comparison. iii) Plots of relative length change of the specimens as a function of time during isothermal tempering.
3.6 Thermodynamical calculations
Thermo-Calc, a thermodynamical calculation software was developed in the mid-70’s. It is now commonly used for development and evaluation of alloys in academic and industrial research. The software was used to calculate the equilibrium composition of different phases and calculations of phase diagram. An internal data base ‘Tooling 11’ specially developed for Uddeholms steels was used when performing the calculations.
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EXPERIMENTAL
3.7 Mechanical testing
3.7.1 Hardness
Hardness measurements were carried on specimens having similar dimensions as those used for microscopy (see Section 3.3). Rockwell C hardness was measured with a Zwick/Roell ZHR hardness tester at room temperature. Each hardness value is reported as an average of five measurements.
3.7.2 Compressive strength
The specimens for uniaxial compression testing had a dumb bell shape with a length of 42 mm and diameter of 16 mm. They were taken from the centre of the steel billet and were tested in longitudinal direction (see fig. 10). The uniaxial compression test was performed with an MTS Universal Testing machine at room temperature. The yield strength was defined at a compression of 0.2 % plastic strain. Each yield strength value is an average of three measurements.
Rolling direction
Charpy V-notched and un-notched Radial direction impact testing specimen in transversal direction
Figure 10 Illustration of specimens taken from the soft annealed bar showing direction of rolling and the directions of testing.
3.7.3 Impact toughness
Specimens for un-notched and V-notched impact toughness testing were produced with the dimension of 55 x 7 x 10 mm and 55 x 10 x 10 mm (l x b x h), respectively (see fig 10). They were machined from the centre of the steel
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EXPERIMENTAL billet in the transversal direction. The notched specimens contained a 2 mm deep 45° V-notch with a radius of 0.25 mm with the notch pointing in the radial direction. The impact testing was carried out at room temperature with an Amsler/Roell RKP 150 equipment. The impact energies are the average of five measurements.
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RESULTS
4 Results
4.1 Mechanical properties
4.1.1 Hardness
The hardness of the specimens in as-quenched condition was for all austenitisation temperatures around 65 HRC. The hardness as a function of tempering temperatures (2x2h, refer Table 6) after austenitisation at temperatures of 1020°C, 1050°C and 1075°C is presented in fig. 11. The hardness showed similar trends for all austenitisation temperatures. Hardness was around 60 HRC when the steel was tempered at 200°C. Tempering at temperatures between 250°C to 450°C, resulted in a lower hardness of about 56-59 HRC. The hardness reached a maximum (secondary hardening peak) when tempering at 500°C for specimens austenitised at temperatures of 1020°C and 1050°C. The specimens austenitised at 1075°C showed the same hardness of 62 HRC when tempered at 500°C or 525°C. It can be noticed that the highest hardness at secondary hardening peak was seen for specimen austenitised at 1075°C and the lowest hardness for the specimens austenitised at 1020°C.
Figure 11 Hardness as a function of tempering temperature (2 x 2h) for specimens that were austenitised at 1020°C, 1050°C or 1075°C. The secondary hardening peak was at 500°C for specimens austenitised at 1020°C and 1050°C while specimens austenitised at 1075°C had similar hardness at 500°C and 525°C. (Paper B).
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RESULTS
Tempering at temperatures above 525°C resulted in lowering of the hardness for all austenitisation temperatures. When comparing hardness of specimen austenitised at different temperatures, it was noticed that the specimens austenitised at 1075°C had the lowest hardness below 500°C and highest hardness above.
4.1.2 Compressive strength
Results from uniaxial compression testing for austenitised and tempered specimens (2x2h, see. Table 6) are presented in Table 8. Notice that, for all austenitisation temperatures, the compressive yield strengths were highest (around 2300 MPa) for the specimens tempered at 525°C and lowest when tempered at 200°C (2192-1852 MPa). Austenitisation at higher temperatures resulted in a lower compressive yield strength when tempering at 200°C.
Table 8 Compressive yield strength for specimens that were austenitised at 1020°C, 1050°C and 1075°C and tempered at 200°C and 525°C.
Austenitisation Tempering Compressive yield temperature temperature strength (˚C) (˚C) (MPa) 200 2192 ± 12 1020 525 2319 ± 9 200 2041 ± 15 1050 525 2337 ± 10 200 1852 ± 33 1075 525 2299 ± 5
4.1.3 Impact toughness
The results from the un-notched impact testing of austenitised and tempered (2x2h, see Table 6) specimens are shown in Table 9. Austenitisation at a higher temperature reduced the impact toughness. It was also noticed that the impact energies for specimens tempered at 200°C were higher compared to impact energies after tempering at 525°C.
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RESULTS
Table 9 Un-notched impact energies of specimens that were austenitised at 1020°C, 1050°C or 1075°C and tempered at 200°C or 525°C.
Austenitisation Tempering Un-notched temperature temperature impact energy (˚C) (˚C) (J) 200 156 ± 20 1020 525 128 ± 20 200 131 ± 18 1050 525 76 ± 14 200 37 ± 5 1075 525 21 ± 2
The results from the Charpy un-notched and V-notched impact testing of the specimens austenitised at 1050°C, tempered for one, two and three times (see Table 7) are presented in fig. 12. The un-notched impact energy of the specimens were lowest (60 J) after the first tempering treatment. There was an increase in the impact energy after the second tempering (75 J). However, the difference in the impact energy after the second and third tempering was very small.
Charpy impact energy of V-notched specimens did not show any significant changes with the tempering treatment. The only prominent difference is that the impact energies showed less scatter after the third tempering.
Figure 12 Charpy un-notched and V-notched impact testing of specimens that were austenitised at 1050°C and tempered one, two or three times at 525°C. The un- notched impact energies after two or three tempering treatments were higher than after tempering one time. (Paper A).
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RESULTS
4.1.3.1 Fractography
The macrographs of fractured surfaces after un-notched impact testing of tempered specimens (2x2h, see table 6) are shown in fig. 13. It was evident from the uneven fracture surfaces of specimens austenitised at 1020°C and 1050°C (see figs. 13 a, b, c and d) that they had absorbed more impact energy than specimens austenitised at 1075°C having flatter fracture surfaces (fig. 13e and f).
Figure 13 Fracture surfaces after un-notched impact testing a), b), c) and d) showed high resistance to fracture while e) and f) showed flatter fracture surface indicating a less ductile fracture.
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RESULTS
Examples of fracture surfaces of un-notched impact specimens with high impact energy and low impact energy are presented in fig. 14. The fracture surfaces of higher impact energy specimens had signs of multiple secondary crack initiations (fig. 14a). However, specimens that absorbed lower impact energies had a flatter surface and with what appeared to be a single point crack initiation.
Impact Impact Secondary crack initiation Crack initiation
Crack propagation
Impact Impact
a) High impact energy (156 J) b) Low impact energy (37 J)
Figure 14 Fracture surfaces of un-notched impact specimens. The fracture that absorbed: (a) high impact energy showed secondary crack initiations and (b) low impact energy showed single point crack initiation.
4.2 Microstructures
4.2.1 Soft annealed microstructure
A SEM (backscattered electrons) micrograph of the investigated steel in soft annealed condition is shown in fig. 15. It shows three different kinds of carbides labelled as grey, bright, and dark.
The results from EDS analysis of carbides in fig. 15 are shown in Table 10. The grey carbide (EDS. 1) is chromium-rich, the bright carbide (EDS. 2) is molybdenum-rich while the dark carbide (EDS. 3) is vanadium-rich. The bright and dark carbides are very small therefore a larger contribution in their composition comes from the matrix.
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RESULTS
Bright carbide, EDS. 2
Dark carbides, EDS. 3
Grey carbides, EDS. 1
10 μm
Figure 15 A SEM (backscattered electrons) micrograph of the investigated steel in soft annealed condition. Arrows shows grey, bright and dark carbides. EDS 1, 2 and 3 refer to analyses in Table 10.
Table 10 SEM-EDS analyses (wt.%) of carbides in the soft annealed microstructure of the investigated steel.
EDS. Cr Mo V Fe + C 1. 28 6 6 Bal. 2. 6 14 1 Bal. 3. 6 7 9 Bal.
4.2.2 As-quenched microstructures
The microstructure of specimens that were austenitised at temperatures of 1020°C, 1050°C or 1075°C are presented in fig. 16. Notice that the steel when austenitised at 1020°C (fig. 16a) showed very fine prior austenite grains. The prior austenite grains were somewhat coarsened when the specimen was austenitised at 1050°C (see fig. 16b). However, the grain size increased drastically after austenitisation at 1075°C (see fig. 16c). The average prior austenite grain sizes when measured according to ASTM standards were 11 μm, 16 μm and 45 μm, after austenitisation at 1020°C, 1050°C, and 1075°C respectively.
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RESULTS
a) 1020°C
50 μm
b) 1050°C
50 μm
c) 1075°C
50 μm
Figure 16 Micrographs showing prior austenite grains after austenitisation at 1020°C, 1050°C, and 1075°C. The prior austenite grains are finer when austenitised at 1020°C and 1050°C while they are drastically coarsened when austenitised at 1075°C.
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RESULTS
a) 1020°C
Undissolved carbide
Blocky light region
Dark region
20 μm
b) 1050°C
Blocky light region
Undissolved carbide
Dark region
Light region (elongated feature) 20 μm
c) 1075°C Dark region
Undissolved carbide
Blocky light region
Light region (elongated feature)
20 μm
Figure 17 Hardened microstructures of the investigated steel after austenitisation at a) 1020°C b) 1050°C and c) 1075°C. All microstructures showed undissolved carbides, dark and light regions. The light regions were either blocky or elongated.
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RESULTS
High magnification optical micrographs of the specimens austenitised at 1020°C, 1050°C or 1075°C are presented in fig. 17a, b and c, respectively. All microstructures showed undissolved carbides along with dark and light etching regions. It can be further noticed that the light regions consisted of blocky as well as long elongated features (prominent in fig. 17c).
A SEM (backscattered electrons) micrograph after austenitisation at 1050°C is shown in fig. 18. Notice that the microstructure after austenitisation shows only grey carbides.
Grey carbide, EDS. 1
10 μm
Figure 18 SEM (backscattered electrons) micrograph after austenitisation at 1050°C. Only grey carbides remain undissolved.
SEM-EDS analysis of the grey carbide revealed that it is chromium-rich as shown in Table 11. It is in agreement with the microstructure of the soft annealed specimen (fig. 15 and Table 10).
Table 11 SEM-EDS analysis of the grey carbide (wt. %) that remained undissolved after austenitisation at 1050°C.
EDS. Cr Mo V Fe + C 1. 26 7 6 Bal.
The volume fraction of retained austenite as measured with XRD increased with austenitisation temperature and was 19 %, 23 % and 28 % in specimens austenitised at 1020°C, 1050°C or 1075°C, respectively.
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RESULTS
4.2.3 Tempered microstructures
4.2.3.1 Single tempering treatments
The as-quenched microstructure of the investigated steel with undissolved carbides, etched regions and regions that were not etched having a blocky appearance is shown in fig. 19a. The microstructure of the specimens tempered at 525°C for holding times of 0.1, 0.5, 1, 2 and 6h (see Table 5) were similar i.e. consisting of similar fractions of etched and less etched blocky regions. A representative microstructure, tempered for 2 h is presented in fig. 19b. It shows undissolved carbides (larger), carbides (light) in etched regions (less than a few μm) and unetched blocky regions. However, the amount of blocky regions in the microstructure was smaller than in the as-quenched microstructure.
a) Q50 b) Tempered at 525°C (2 h)
Etched region Blocky unetched region
Blocky unetched region Etched region
Undissolved carbide Carbides
5 μm 5 μm
Figure 19 SEM micrograph of a) an as-quenched specimen austenitised at 1050°C and b) after tempering for 2 hours. Both microstructures showed undissolved carbides, together with etched and unetched regions. A difference is the small carbides in etched regions that were present only in the tempered microstructure.
The microstructure of a specimen that was tempered for 10 h is shown in fig. 20a. It shows carbides (light) in regions that were etched and some blocky regions which were less etched. Notice that the blocky regions are not prominent as in the microstructures for short aging time (compare figs. 19 and 20). A higher magnification micrograph of the same region in fig. 20b revealed that the blocky regions contain very small carbides.
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RESULTS
a) Tempered at 525°C (10 h) b) Tempered at 525°C (10 h)
Blocky region Carbides in etched region
Blocky region Very small carbides
5 μm 500 nm
Figure 20 SEM micrographs of a specimen austenitised at 1050°C and tempered at 525°C for 10 hours. a) Etched regions with carbides (light) and less etched blocky regions appear. b) The blocky regions contained very small carbides as revealed at a higher magnification.
The specimens that were austenitised at 1050°C were tempered with varying holding time at 525°C (see Table 5). The fractions of retained austenite were measured and are presented in Table 12. The retained austenite fraction decreased when the holding time at 525°C was increased. Tempering for 6 h or 10 h at 525°C resulted in less than 2 % retained austenite.
The lattice parameter of retained austenite was measured and is also presented in Table 12. The lattice parameter increased for the shortest holding time of 0.1 h when compared with the lattice parameter for the as-quenched specimen. Afterwards, the lattice parameter decreased with longer holding times.
Table 12 Fractions and lattice parameters of retained austenite when tempered at
525°C with increasing holding times or at 600°C for 30 h (T30).
Specimen Holding time RA Lattice parameter (aȖ) (h) (vol. %) Å
Q50 - 24 ± 2 3.6037 ± 0.0005
T0.1 0.1 22 ± 2 3.6070 ± 0.0005
T0.5 0.5 17 ± 2 3.6032 ± 0.0002
T1 1 12 ± 2 3.6020 ± 0.0008
T2 2 5 ± 2 3.5998 ± 0.0006
T6 6 <2 -
T10 10 <2 -
T30 30 <2 -
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RESULTS
The specimen that was austenitised at 1050°C and tempered at 600°C for 30 h contained less than 2 % retained austenite. The microstructure of this specimen is shown in fig. 21. It can be noticed that the microstructure contains undissolved carbides, etched areas with dense precipitation of carbides and regions which appear to be blocky. A higher magnification micrograph is presented in fig 21b showing that also the blocky regions contain carbides. The carbides in the more heavily etched regions appear light while the carbides in blocky regions are grey.
a) Tempered at 600°C (30 h) b) Tempered at 600°C (30 h) Dense precipitation of carbides
Carbides (light)
Blocky region
Carbides (grey)
5 μm 1 μm Undissolved carbides
Figure 21 SEM micrograph of specimen austenitised at 1050°C and tempered at 600°C for 30 hours. a) The micrograph shows undissolved carbides, dense precipitation of carbides in etched areas and regions having blocky appearance. b) A higher magnification reveals that the blocky regions contain carbides (grey).
4.2.3.2 Double tempering at 200°C to 600°C
Optical micrographs of specimens austenitised at 1075°C and tempered (2x2h, refer. Table 6) at 200°C or 525°C are shown in figs. 22a and b, respectively. The microstructure after tempering at 200°C showed undissolved carbides, blocky light regions and dark regions. However, there is a prominent difference in the microstructure when the steel was tempered at 525°C as it does not contain any blocky light regions. Instead it shows a dark needle like structure.
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RESULTS
a) 1075°C, 200°C (2x2h) b) 1075°C, 525°C (2x2h)
Undissolved carbide
Blocky light region
Dark region 20μm 20μm
Figure 22 Microstructures (SEM micrographs) of specimen tempered 2x2h at a) 200°C or b) 525°C. Notice that the blocky light regions are present after tempering at 200°C but not after tempering at 525°C.
A higher magnification SEM micrograph of the specimen tempered at 525°C is presented in fig. 23. It shows that the martensitic microstructure contained fine carbides (light).
1075°C, 525°C (2x2h)
5 μm
Figure 23 A SEM micrograph of the specimen austenitised at 1050°C and tempered at 525°C (2x2h). There are small carbides (light) in a tempered martensitic matrix.
The amount of retained austenite as a function of tempering temperature (2x2h) for specimens austenitised at 1020°C, 1050°C or 1075°C is presented in fig. 24. The amounts did not change much for low tempering temperatures (200-450°C) where retained austenite content was fairly stable. When tempering between (450-500°C) there was a reduction in the amount of retained austenite and tempering above 500°C resulted in less than 2% retained austenite in the microstructure. This region is marked with a rectangle in fig. 24.
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RESULTS
Figure 24 Fraction of retained austenite as a function of tempering temperature (2x2h) for specimens that were austenitised at 1020°C, 1050°C, or 1075°C. Notice that the fraction of retained austenite becomes less than 2% when tempering above 500°C.
4.2.3.3 Multiple tempering treatments at 525°C
The microstructure of the specimen austenitised at 1050°C and tempered one time at 525°C is presented in fig. 25a. The micrograph shows undissolved carbides, precipitation of carbides in etched regions and unetched blocky regions. The microstructure after the second tempering treatment is presented in fig. 25b. Notice that the microstructure is largely etched showing carbides. However there are a few regions that remained unetched and appeared blocky as labelled in fig. 25b. The microstructure after the third tempering treatment is presented in fig. 25c. There are carbides in etched areas and no blocky regions.
The retained austenite content after austenitisation at 1020°C, 1050°C or 1075°C and tempering one, two, and three times at 525°C (see Table 7) is presented in Table 13. The first tempering treatment resulted in a significant reduction of retained austenite. For example the as-quenched microstructure after austenitisation at 1050°C had 23% retained austenite which was lowered to 7% after the first tempering. After a second or a third tempering treatment (only for 1050°C) less than 2% retained austenite remained.
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RESULTS
a) T50x1
Blocky light regions
Carbides
Undissolved carbides
5 μm
b) T50x2
Carbides
Blocky regions 5 μm
c) T50x3
Carbides
5 μm
Figure 25 Microstructures after a) first tempering, b) second tempering and c) third tempering treatment at 525°C. All microstructures contain carbides in etched areas. Notice that the fraction of blocky regions decreased significantly after the second and third tempering treatments.
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RESULTS
Table 13 Retained austenite for the specimens that were austenitised at 1020°C, 1050°C or 1075°C after single, double and triple tempering treatment at 525°C.
First tempering Second tempering Third tempering Sample Fraction of RA Sample Fraction of RA Sample Fraction of RA (vol. %) (vol. %) (vol. %)
T20x1 <2 T20x2 <2
T50x1 7 T50x2 <2 T50x3 <2
T75x1 12 T75x2 <2
4.3 Transformation behaviour
4.3.1 Hardening
Dilatometry data was gathered during quenching of the specimens that were austenitised at temperatures of 1020°C, 1050°C or 1075°C (see Table 7). The length change of the dilatometry specimens during cooling from 800°C to room temperature is shown in fig. 26a. It also shows an example of the derivative of length change as a function of temperature for the specimen austenitised at 1020°C (black line). Notice that the length of the specimens decreased linearly until 250°C. Below that temperature, an increase in the length of the specimen was observed and bainite (Bs) and martensite start (Ms) temperatures were identified with the help of the derivate of length change.
The relative length change (%) as a function of specimen temperature from 200°C to room temperatures is presented in fig. 26b. Bainite (Bs) and martensite (Ms) transformation start temperatures during cooling for different austenitisation temperatures are marked with arrows in fig. 26b (the method of measurement for bainite and martensite start temperatures is described in section 3.5). Notice that the bainite and martensite start temperatures were lower for specimens that were austenitised at higher temperatures. These results are also presented in Table 14.
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RESULTS
a)
b)
Figure 26 a) Length change as a function of specimen temperature during cooling from the austenitisation temperatures of 1020°C, 1050°C or 1075°C. Notice identification of martensite and bainite start temperatures on the derivative of length change (black line). b) Relative length change (%) as a function of temperature from 200°C to room temperature. The martensite and bainite start temperatures were lower when austenitisation took place at higher temperatures.
Table 14 Martensite and bainite start temperatures for specimens austenitised at 1020°C, 1050°C or 1075°C.
Martensite Bainite Sample start temperature start temperature Ms (˚C) Bs (˚C)
1. Q20 180 230
2. Q50 150 200
3. Q75 140 190
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RESULTS
4.3.2 Single tempering
The tempering behaviour during the first tempering at 525°C of specimens austenitised at temperatures of 1020°C, 1050°C, or 1075°C (see Table 7) is presented in fig. 27. It shows the change in relative length of the specimen as a function of time for 2 hours. Notice that the specimens initially showed expansion for approximately 5 minutes followed by contraction during the rest of the tempering time. Contractions were measured as a difference from the point of maximum expansion (i.e after 5 min of tempering) to the point of minimum expansion of the specimen (i.e after 2 hours of tempering) as shown in fig. 27. The specimens austenitised at 1020°C or 1050°C showed a contraction of 0.007% during tempering while the specimen austenitised at 1075°C showed a contraction of approximately 0.006%.
Figure 27 Relative length change (%) of the specimens as a function of time (2h) during the first tempering at 525°C. Specimens that were austenitised at 1020°C, 1050°C had a contraction of 0.007% while specimen austenitised at 1075°C showed a contraction of approximately 0.006%.
When cooling to ambient temperature after the first tempering retained austenite transformation occurred for all specimens. The transformation temperatures were evaluated and are presented in Table 15. It can be noticed that a higher austenitisation temperature resulted in a lower Ms temperature.
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RESULTS
Table 15 The martensite start temperatures during cooling after a single tempering treatment.
Sample Martensite start temperature Ms (˚C)
T20x1 190
T50x1 170
T75x1 160
The relative length change as a function of time for a specimen that was austenitised at 1050°C and tempered at 525°C for 10 h (see Table 5) is presented in fig. 28. Notice that the total contraction of the specimen during 10 h was 0.01 %.
Figure 28 Relative length change (%) as a function of tempering time for specimens austenitised at 1050°C and when tempering at 525°C for 10 h or at 600°C for 30 h. A contraction occurred when tempering at 525°C with a total contraction of 0.01%. However, when tempering at 600°C, a contraction of 0.01 % occurred during the first 2 hours followed by 0.1% expansion until the end of tempering.
Relative length change as a function of time for a specimen that was austenitised at 1050°C and tempered at 600°C for 30 h (refer Table 5) is also shown in fig. 28. The specimen contracted about 0.01% during approximately the first 2h. Afterwards, the specimen expanded until the end of tempering. The expansion was 0.1%.
Martensite start and finish temperatures were estimated and are provided in Table 16 for specimens austenitised at 1050°C and tempered for holding times of 0.1, 0.5, 1, 2, 6 or 10 hours (Table 5) at 525°C. Increasing the holding times at
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RESULTS
525°C resulted in higher martensite start temperature as presented in fig. 29. The specimens tempered for longer holding times of 6 or 10 h also showed martensite finish temperatures. The finish temperatures were also found to increase with longer holding times. No martensite start temperature was found for the specimen tempered for the shortest holding time (0.1 h).
Figure 29 Length change as a function of specimen temperature during cooling to room temperature after tempering at 525°C for 0.1 (T0.1), 0.5 (T0.5), 1 (T1), 2 (T2), 6
(T6) and 10 hours (T10) and specimen tempered at 600°C for 30 hours (T30). Martensite start temperature increased with longer time at holding temperature. Notice that no martensitic transformation was observed for the specimen tempered at 525°C for 0.1 h or for the specimen tempered at 600°C for 30 hours.
For the specimen austenitised at 1050°C and tempered at 600°C for 30 h, no martensitic transformation was observed during cooling. This suggests that the retained austenite was transformed during tempering.
Table 16 Martensite start and finish temperatures measured during cooling after tempering at 525°C for varying holding times and at 600°C for 30 h.
Specimen Holding time (h) Ms Mf
T0.1 0.1 - -
T0.5 0.5 40 -
Tempering at 525°C T1 1 120 -
T2 2 170 -
T6 6 240 140
T10 10 280 180
Tempering at 600°C T30 30 - -
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RESULTS
4.3.3 Double tempering
The double tempering treatments of the specimen that were austenitised at 1020°C, 1050°C or 1075°C (Table 5) resulted in contraction as shown in fig. 30. The specimen austenitised at 1020°C had a contraction of approximately 0.001 %. The specimen that was austenitised at 1050°C showed a contraction of approximately 0.002 % while the specimen austenitised at 1075°C had a contraction of approximately 0.003 %. These contractions are very small compared to the contraction occurring during the single tempering treatment (compare with fig. 27).
Figure 30 Relative length change (%) of specimens during the second tempering treatment at 525°C. Notice that the contraction was lowest for specimen austenitised at 1020°C and highest for the specimen austenitised at 1075°C.
During cooling from the second tempering, there was retained austenite transformation only for specimens that were austenitised at 1050°C or 1075°C. The Ms temperatures are presented in Table 17.
Table 17 The martensite start temperatures during cooling after the second tempering treatments.
Second tempering Sample Fraction Martensite of RA start temperature (vol. %) Ms (˚C)
T20x2 <2 -
T50x2 <2 180
T75x2 <2 180
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RESULTS
4.4 Thermodynamic calculations
The phase diagram for the typical composition of the investigated steel with varying carbon content for the temperatures from 800°C to 1500°C is shown in fig. 31. The carbon content of the investigated steel is marked with a dashed line. The marked points in the figure shows the critical range of temperatures from 980°C to 1080°C where austenite and M7C3 are stable phases. However, above 1080°C austenite is the only stable phase.
Figure 31 Equilibrium phase diagram of the investigated steel as a function of varying carbon content. The dashed line represents the carbon content of the investigated steel.
The calculated equilibrium composition of austenite for the composition of the investigated steel at temperatures of 1020°C, 1050°C or 1075°C is presented in Table 18. Notice that austenitisation at higher temperatures dissolves more carbide forming elements (C, Cr, Mo, and V) in austenite.
Table 18 Calculated chemical composition of austenite (wt.%) at temperatures of 1020°C, 1050°C or 1075°C.
Austenitisation C Cr Mo V Si Mn Fe temperature (°C) 1020 0.54 4.34 2.10 0.33 0.20 0.50 Bal. 1050 0.61 4.62 2.23 0.40 0.20 0.50 Bal. 1075 0.67 4.80 2.28 0.46 0.20 0.50 Bal.
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RESULTS
The calculated chemical composition of equilibrium carbides at 860°C is provided in Table 19. This calculation was performed to permit a comparison with the carbides in soft annealed condition. Notice that M7C3 carbides are chromium-rich, M6C carbides are molybdenum-rich while MC carbides are vanadium-rich.
Table 19 Calculated chemical composition (wt.%) of carbides for the investigated steel at the temperature of 860°C.
Carbides C Cr Mo V Mn Fe
M7C3 8.42 41.11 9.52 7.22 0.38 Bal.
M6C 2.66 3.31 57.75 1.30 - Bal. MC 15.07 6.25 20.62 55.43 0.02 Bal.
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DISCUSSION
5 Discussion
The main aim of this thesis has been to try to understand the effect of different heat treatments on the microstructure and properties of a 5 wt.% Cr cold work tool steel. Thus, the work has included both different austenitisation temperatures and various tempering treatments. Some mechanical tests have been performed as well, but most of the results rely on microstructural examination.
5.1 Microstructure
The first question to discuss is the choice of austenitisation temperature. From the size of prior austenite grains (figure 16) it is clear that austenitisation at 1075°C caused a very large growth of austenite grains (from 11 to 45 µm). The very large prior austenite grains also caused a very coarse martensite structure (see fig. 17) after cooling to room temperature, which is deemed not desirable to have in a tool steel. Thus, from these results, austenitising at 1020°C or 1050°C seems to be the only realistic alternatives.
The as-quenched microstructures revealed a mixture of dark and light regions for all specimens (see figs. 17 or 19). The dark regions seemed to be more interconnected, while the light regions consisted of smaller separate “islands”, with a more angular shape (fig 19a). As seen in fig 19b, after heat treatment at 525°C for 2 hours, the dark regions have become more frequent and the angular light regions have become less, but far from being completely dissolved. It can also be noted that a large number of fine precipitates can be seen in the dark regions. With increased tempering time at 525°C, up to 10 hours, a further decrease in the amount of light angular islands is seen and also, as shown in fig 20b, there seems to be precipitation in the angular regions. However, when double tempering is performed, i.e. 2+2 hours at 525°C, the angular features in the microstructure are more or less gone depending on the austenitising temperature (fig. 25).
It is of course tempting to assume that the overall matrix in the microstructure in as-quenched condition is martensite (possibly with some bainite) while the angular light regions are retained austenite. However, just a coarse evaluation of the volume fractions of the different constituents show that there is much more angular light regions than the amount of retained austenite as measured by XRD. Thus, clearly the microstructure needs to be studied in much more detail, by for example transmission electron microscopy, before any more fundamental
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DISCUSSION understanding of the microstructure in the as-quenched condition can be obtained.
Literature survey suggests that there are not many published articles on 5 wt.% Cr cold work tool steel. It is largely accepted that the carbon content of Fe-C alloys has a strong effect on the morphology of as-quenched martensite, for instance there is a morphological change of the martensite from lath martensite at low carbon contents to plate martensite at high carbon contents [23]. For medium carbon steels (as the investigated steel) a mix of plate and lath martensite exists [7, 24]. A study for a 12 and 8 wt.% Cr steel [25] suggested a strong effect of austenitisation temperature on the morphology of martensite. For example a higher austenitisation temperature with more dissolved alloy content produces plate martensite rather than lath martensite. This is reflected in the microstructures in fig. 17. Moreover, the etching of martensite depends on both its orientation and morphology [26]. Thus, it seems plausible to believe that some of the fresh martensite may not etch and can appear blocky.
The direct observation of retained austenite in the microstructure has always been difficult particularly if the retained austenite is available in low concentrations. It is known that the retained austenite present in low carbon steels have a thin film structure [25]. It is usually found in between the laths of martensite or bainite [25, 27]. However for steels containing higher carbon, retained austenite can also be present in the form of islands. Yaso et.al. tried to characterise the blocky retained austenite by EBSD and thin film retained austenite by TEM for 12 wt.% Cr and 8 wt.% Cr alloy steels [28]. The results showed that one technique alone is not enough to fully quantify the retained austenite content when compared with XRD results. Due to the given facts, several authors referred the blocky regions as M/A blocks [25, 28, 29] although is more frequently used for ferritic bainitic steels. Thus, in the light of published literature, the interpretation of the microstructures presented in this study seems to agree with of what already has been published.
5.2 Effects of tempering
Tempering is the final heat treatment on the tool. Therefore it is important that the tool attains an appropriate microstructure and good mechanical properties. One important aspect is the presence of retained austenite in the as-quenched microstructure. With tempering treatments and the associated cooling the steel is transformed so that retained austenite is avoided in the final microstructures. It is well known that retained austenite in hardened steels can transform during
64
DISCUSSION tempering treatments or the associated cooling. The temperatures at which retained austenite starts to transform is dependent on the steel composition.
It was seen that retained austenite transformations occurred in steels tempered above 500°C (fig. 24). The variations of holding time during the single tempering treatment at 525°C will be discussed further to understand how retained austenite transformation occurred.
5.2.1 Retained austenite transformation to martensite
5.2.1.1 Single tempering treatments
The specimens after austenitisation at 1020°C, 1050°C or 1075°C contained various proportions of un-tempered martensite (and possibly bainite) and retained austenite (in total 100 %) resulting in a comparatively large contraction during the first tempering at 525°C (fig. 27). The contraction of the specimen during tempering can be explained by precipitation of carbides in the martensite (possibly bainite) and retained austenite [30, 31]. Unlike martensite, bainite contains only a slight excess of carbon in solution thus a large contribution from bainite is not expected during tempering [16]. However, cooling from the first tempering caused partial transformation of retained austenite to martensite (see Ms in table 15) at least for specimens austenitised at 1050°C or 1075°C. For example the specimen austenitised at 1050°C, 24% retained austenite was reduced to 7 % after first tempering. The microstructure tempered for 2 h however, showed quite a lot of blocky regions (fig. 19b). Thus, it was estimated that the blocky regions were consisting of a mix of fresh martensite (17%) and retained austenite (7 %).
Dilatometry experiments for single tempering treatments at 525°C with variation in holding time also showed contraction of specimens during the holding time (see section 4.3.2 or paper A or C). This is already explained as precipitation of carbides occurred in the martensite (possibly bainite) and retained austenite [30, 31]. Microscopy indeed showed precipitation of carbides in martensite (or bainite) in the microstructure of specimen tempered for 2 hours (fig. 19). Direct observation of the precipitation of carbides in retained austenite was not possible because the precipitates were too small to be revealed by SEM in the microstructure for short holding times. However, it can be concluded from literature that precipitation in retained austenite can occur fairly quickly. For example, for a hot work tool steel, it was reported [32] that the formation of carbides in retained austenite start already when heating to the tempering temperature. On the other hand Kulmberg reported [33] for a high speed steel that carbides form in retained austenite during tempering which on
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DISCUSSION cooling transforms to martensite. Most of the published articles claim the precipitates to be cementite [30, 34, 35] while reference [33] reports precipitation of alloy carbides. Which carbide will form naturally depends on the composition of the steels investigated. Consequently precipitation of alloy carbides can be expected in a tool steel as reported in [33]. The precipitation of carbides progresses with increasing tempering time and makes the retained austenite sufficiently unstable for the martensite transformation to occur on cooling for specimens with holding times from 0.5 to 10 hours.
It was found that longer holding times resulted in a higher Ms temperature (Table 16 or fig. 29) and a lower lattice parameter for the retained austenite (Table 12). This is in line with earlier observations as reported for example in [33] for Ms temperature and in [36] for the lattice parameter. This provides an opportunity to further investigate the changes in the composition of the retained austenite occurring during tempering.
Equation 1 [37] was used to estimate the carbon content of the retained austenite by using measured Ms temperatures (Table 16). The equation was selected because of its validity for steels containing 0.6 wt.% C and 5 wt.% Cr, which is similar to the composition of the present steel even though the equation does not include a factor for vanadium content.
Ms = 539 - 423 C – 30.4 Mn – 17.7 Ni – 12.1 Cr – 11.0 Si – 7.0 Mo Eq.1
The earlier relationships developed to determine the lattice parameter of the retained austenite and its carbon content [38, 39] did not include the effect of substitutional alloying elements. However, a modified equation which includes the main alloying elements of the studied steel is presented as Eq. 2 [40] and was used to calculate the carbon content of retained austenite from the measured lattice parameters (Table 12). The austenite composition used (except for C) when applying eq.1 and eq.2 was calculated by Thermo-Calc (see Table 18). aγ = 3.573 + 0.033 Cγ + 0.00095 Mn – 0.0002 Ni + 0.0006 Cr + 0.0031 Mo + 0.0018 V Eq.2
The carbon content of retained austenite calculated with Thermo-Calc and Eq. 1 and Eq. 2 was plotted against holding time at 525°C and is presented in fig. 32.
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DISCUSSION
Figure 32 Calculated carbon content of retained austenite (wt.%) as a function of holding time at 525°C. The carbon content was calculated with Thermo-Calc, Eq. 1 (Ms) and Eq. 2 (lattice parameter). The calculated carbon content decreased with longer holding times.
The trends are very similar for the two curves showing an initial increase (see paper C for a detailed discussion on the initial increase) followed by a decrease for longer holding times. It can be noticed that the Thermo-Calc calculation (0.61 % C, Table 18) is between the two curves and very close to the as- quenched value calculated from lattice parameter measurements (0.59 % C, fig. 32). Although there is some difference between the two curves it was concluded that the trends rather than absolute values are important for understanding the retained austenite transformation. Hence, the increase in Ms temperature and the decrease in lattice parameter for longer aging times is associated with depletion of carbon in retained austenite due to precipitation of carbides. This conclusion is supported by the observation of precipitation of very small carbides in the blocky regions in the microstructure after tempering for 10 h (fig 20).
5.2.1.2 Multiple tempering treatments at 525°C
The primary aim of studying multiple tempering treatments using dilatometry was to understand the transformation of retained austenite and the tempering of fresh martensite in the subsequent tempering treatments. It is known that fresh martensite is brittle and can be detrimental for toughness and therefore should be avoided in the final microstructure.
In order to completely understand the transformation of retained austenite and tempering of fresh martensite, it was necessary to estimate the fraction of fresh
67
DISCUSSION martensite for the microstructures after the first, second and third tempering treatments. Therefore by using the XRD results (Table 13) the fraction of fresh martensite was estimated and is presented in Table 20 (see paper A for details on the calculation of fresh martensite).
Table 20 Fraction of fresh martensite (FM) and bainite (B) after the first, second and third tempering treatments.
As-quenched First tempering Second tempering Third tempering FM + B FM FM FM Sample Sample Sample Sample (vol. %) (vol. %) (vol. %) (vol. %)
Q20 81 T20x1 19 T20x2 -
Q50 76 T50x1 17 T50x2 7 T50x3 -
Q75 72 T75x1 16 T75x2 12
The contraction of dilatometry specimen during the first tempering and subsequent retained austenite transformation on cooling has already been discussed in section 5.2.1. During the second tempering the contractions were relatively smaller (from 0.001% to 0.003 %) than the first tempering (from 0.006% to 0.007%). This was due to the presence of lower volume fractions of retained austenite and fresh martensite in the specimen tempered once i.e. for specimen austenitised at 1050°C, 17% fresh martensite and 7% retained austenite was present, see Table 20. A similar result can be found in published data of [30] showing that a fully martensitic steel has a larger contraction during tempering than a steel containing a mix of bainite and retained austenite.
A higher austenitisation temperature also resulted in a larger contraction during the second tempering than a lower austenitisation temperature due to the total higher contents of fresh martensite and retained austenite. For example the contraction was 0.003 % for the highest austenitisation temperature with 16 % FM + 12% RA (total 28%) compared to 0.001 % contraction for the lowest austenitisation temperature with 19 % FM and no retained austenite. What proportion of the contraction can be attributed to fresh martensite, tempered martensite and retained austenite cannot be estimated from the present study.
It was seen that the Ms temperature for the specimen austenitised at 1075°C was higher during cooling after the second tempering (180°C) when compared with the first tempering treatment (160°C, compare Tables 15 and 17). This is a consequence of carbon depletion in retained austenite as discussed in 5.2.1.1. After the third tempering, XRD results showed no retained austenite and
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DISCUSSION martensitic transformation was not observed during cooling i.e. the microstructure was completely tempered.
The above discussion was aimed to increase the understanding of the tempering process. Related studies reported in literature have [1, 41] mainly been conducted on high speed steels which contain a high content of alloying elements resulting in a high content of retained austenite after quenching.
5.2.2 Retained austenite transformation to ferrite and carbides
The tempering treatment at 600°C resulted in contraction of the specimen for the first two hours (see fig. 28). This is interpreted as an effect of the precipitation of carbides in martensite and retained austenite as discussed in section 5.2.1.1. However, after the initial contraction the specimens expanded. Literature suggests that the expansion of specimen during tempering is due to diffusional transformation of retained austenite to ferrite and carbides [29, 30, 33]. The microstructure of the specimen certainly showed blocky regions containing carbides (grey) in ferrite (fig. 21). Further observations supporting this are that no martensitic transformation was observed during cooling (Table 16) and that no retained austenite was present after the tempering treatment (Table 12). Thus it was concluded that all the retained austenite was transformed during holding at the tempering temperature.
Similar experiments on a variety of steels can be found in literature [29, 30, 31, 33, 42]. Relevant examples are for a 5 wt. % Cr hot work tool steel containing 0.37 wt.% C during tempering at 610°C [33] while Podder and Bhadeshia studied tempering at 450°C of a bainitic steel with 0.22 wt.% C - 3 wt.% Mn - 2.03 wt.% Si [30]. References [33] and [30] showed a contraction for the first 20 minutes and 1 hour, respectively, followed by expansion. The rate of carbide precipitation is related to the carbon diffusivity in retained austenite as reported in [43]. Thus, it is not unexpected that the process is more rapid at 610°C than at 450°C. It can be noticed that the bainitic steel does not include carbide forming alloy elements which may affect the precipitation of carbides at high temperatures.
The hot work tool steel in reference [33] has about half the carbon content of the cold work steel used in the present study (Uddeholm Caldie). It can be expected that a higher carbon and alloying content of the retained austenite will delay the transformation of retained austenite into ferrite and cementite. This can explain that the transformation was not observed until after 2 hours in
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DISCUSSION
Caldie but was seen already after 20 min in the 5 wt. % Cr hot work tool steel with lower carbon content.
There are different opinions about the mechanism of retained austenite transformation to ferrite and carbide. Van Genderen et. al. [34] divided that the retained austenite decomposition of a FeC model alloy into two successive ways. A preceding ferrite formation occurs before the final transformation to ferrite and cementite take place. However, the results presented in this thesis follow well reference [33] and can be described as follows. The retained austenite is saturated with alloying elements and is thermodynamically unstable. Thus the precipitation of carbides from retained austenite is the first process depleting carbon from retained austenite and permits a second process where ferrite forms. However ferrite has a low solubility for carbon and promotes carbide precipitation simultaneously.
5.2.3 Mechanical properties
5.2.3.1 Hardness and compressive strength
The high hardness of as-quenched specimens was due to the hard and brittle nature of un-tempered martensite. During tempering, carbides precipitated in the martensite and a significant reduction in the hardness was seen for all tempered specimen. For example, the reduction in hardness was 5 HRC when the steel was tempered at 200°C while the reduction was 4 HRC (for specimen austenitised at 1020°C) and 3 HRC (for specimen austenitised at 1050 or 1075°C) after tempering at 500°C. It is known that the precipitation of cementite occurs in steels when tempering from 200°C to 350°C. The coarsening of cementite and the effects on the hardness of the steel is well described by Speech and Leslie [7]. The change in hardness of the present steel when tempered at temperatures between 200°C to 475°C shows a decrease and a regain in hardness. It can be explained as formation of array of cementite rods at temperatures above 250°C to 350°C while the loss of rod shape and gradual spherodisation of cementite at temperatures above 350°C helps in regaining the hardness.
It is known that the secondary hardening peak occurs due to the precipitation of MC and M2C carbides [17, 18]. The hardness reaches its maximum when the tempering carbides are very small and finely distributed. However, the hardness was lower when tempering above the secondary hardening temperatures (i.e above 525°C). This behaviour is described being due to the formation of coarser
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DISCUSSION secondary particles with a lower density meaning that the average distance between the particles increases and contribute less to the hardness.
The specimens austenitised at 1075°C dissolved a larger amount of alloying elements in the matrix (refer Table 18) resulting in the precipitation of a larger volume fraction of MC and M2C carbides compared to when austenitising at 1020°C or 1050°C. Thus, a higher secondary hardness was obtained for the specimens austenitised at 1075°C. These specimens also contained a higher amount of retained austenite and showed a lower hardness at temperatures below the secondary hardening peak. However tempering above 500°C, resulted in less than 2 % retained austenite and the highest hardness.
The lower compressive strength when tempering at 200°C was due to the presence of a significant amount of retained austenite in these specimens. On the contrary, all austenitised specimens when tempered at 525°C showed higher compressive strength. The increase in compressive strength was due to precipitation strengthening and lower amounts of retained austenite.
5.2.3.2 Impact toughness
The strength and toughness of a steel improves with decrease in grain size [44, 45] following the Hall-Petch relation. Austenitisation at 1020°C and 1050°C for the investigated steel resulted in finer martensite due to fine prior austenite grains. Literature suggests that the martensite blocks and packets are considered effective grains contributing to the strength [46] and toughness [47] of a steel. Most of the studies have shown that the packet size and block width of lath martensite increases with increasing prior austenite grain size [46, 47, 48]. This suggests that a microstructural refinement of martensite increases toughness which is in agreement with observations for specimens austenitised at 1020°C and 1050°C (see fig. 16 and Table 9). Austenitisation at 1075°C on the other hand led to a dramatic coarsening of prior austenite grains which resulted in a coarse martensite structure [24, 49, 50]. However, the trend of reduced toughness with higher austenitisation temperature of the tempered specimens (Table 9) was believed to be due to coarsening of prior austenite grains.
Ductile materials show plastic deformation in the form of necking and formation of shear lips before fracture. Brittle fracture on the other hand happens without being preceded by any plastic deformation. For the investigated steel there were very small signs of plastic deformation on all fractured specimens. Therefore, the steel was considered largely brittle. However, to understand the difference between uneven fracture surfaces that absorbed higher energies and flatter surfaces that absorbed lower energies (see
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DISCUSSION figs. 13 and 14), it is helpful to analyse the material behaviour during un-notched impact testing as illustrated in fig. 33. The fracture initiation is shown in figs. 33a and c while figs. 33b and d show the impact specimens after fracture. For specimens that absorbed higher impact energies, there were multiple cracks propagating resulting in parts of the specimen breaking off during impact testing (see fig. 33b) and the result was a higher material loss than the specimens that absorbed lower impact energies. This was taken as a sign of less brittleness whereas the specimens which had absorbed lower impact energies (see fig. 33d) resulted in small material loss due to what appeared to be single point crack initiation.
Thus, it was concluded that the specimens having coarse martensite basically fractures from a single point crack initiation and show less resistance to crack propagation resulting in a more brittle behaviour. Another contributing factor is a higher alloy content of the martensite which produces a more brittle martensite after austenitisation at 1075°C than at 1020°C or 1050°C.
a) Crack initiation
T 1
Impact
b) Higher impact energy
T f
c) Crack initiation
T 1
Impact d) Lower impact energy
T f
Figure 33 Illustration of fracture during un-notched impact testing. a) Multiple point crack initiation leading to b) a higher material loss. c) Single point crack initiation leading to d) a lower material loss.
72
DISCUSSION
Literature suggests that the presence of retained austenite is beneficial for toughness [51] and provides an increased ductility [16, 52]. The main reason is that the stable retained austenite can postpone crack nucleation and hinder crack propagation at the interface between martensite and carbides [53]. This was believed to be the reason for having a high impact energy for the specimens tempered at 200°C, and agree with the results reported in [12]. However, when tempering at 525°C the precipitation of secondary carbides M2C and MC was abundant in martensite and the formation of fresh martensite occurred during cooling. Thus, it was believed that the high hardness at 525°C was the main reason for reduction in the toughness.
5.3 Heat treatment recommendations
The standard heat treatment procedure practised for the studied 5 wt.% Cr cold work tool steel is a double tempering treatment (2x2 h, see section 3.2.2). Therefore, the aim of this chapter is to suggest ways for improving the heat treatment procedures by investigating the relationship between austenitising temperature, tempering treatment, microstructure and mechanical properties.
The highest austenitisation temperature investigated (1075°C) caused significant coarsening of prior austenite grains and resulted in low impact toughness and would therefore not be recommended. However, for the selection of tempering temperatures it is necessary to discuss briefly the main findings of this work. Tempering treatments for cold work tool steels serve two main purposes depending on the application area of the tool. Firstly, to transform retained austenite and temper martensite and secondly, to achieve better mechanical properties of the tools.
After tempering at 200°C there was a significant fraction of retained austenite. Retained austenite is due to its metastable nature considered detrimental for tool steels as it may transform to martensite on application of load [54]. However, for tool steels this transformation behaviour may lead to dimensional changes which are unacceptable as high dimensional precision is required for finished products.
On the other hand, tempering at 525°C was shown to be a solution to minimize retained austenite, getting maximum hardness and compressive strength while having adequate toughness. It is clear that retained austenite transformed to martensite on cooling after the first and second tempering and thus alternative heat treatment parameters can be suggested. A longer holding time during the first tempering may increase the chance of transforming all retained austenite
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DISCUSSION during the first cooling (see table. 16, T6). Therefore, a second tempering with a short holding time can be used to temper the freshly formed martensite. It was also seen that fresh martensite remained in the microstructure after the double tempering of specimens austenitised at 1050°C or 1075°C. Thus, a third tempering with shorter holding times may be useful to achieve a fully tempered microstructure, in particular when austenitising at high temperatures.
Furthermore, single tempering treatment at 600°C with longer holding time may transform the retained austenite to ferrite and carbides but is not likely to produce the desired hardness. Naturally effects on toughness and compressive strength would have to be explored before recommending any alternative heat treatment procedures.
As a final note, it was concluded that austenitisation at 1020°C or 1050°C of Caldie followed by tempering at 525°C (2x2h) showed sufficient hardness, good compressive strength and adequate toughness and should be a suitable choice for heat treatment of cold work tools (see paper B). Based on these results heat treatment for Uddeholm Caldie has been improved such that austenitisation at 1050°C and a triple tempering treatment has been added in the recommendations.
5.4 Comments on research questions
It was attempted to answer the research questions (see section 1.3) by producing three articles (Paper A, B and C). Paper B deals with research questions 1 and 2:
1. What are the effects of austenitisation temperature on microstructure and hardness? 2. How does the tempering temperature affect the microstructure and mechanical properties?
In this paper the effects of austenitisation and tempering temperature on microstructure and mechanical properties have been discussed. The main findings were that a higher austenitisation temperature will produce a higher secondary hardness. However a too high austenitisation temperature will cause significant coarsening of prior austenite grains which reduce the impact toughness of the material. Therefore, a good combination of properties was obtained when austenitisation occurred at 1020°C or 1050°C followed by tempering at 525°C.
Research question 3 has been discussed in paper A:
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DISCUSSION
3. What are the effects of multiple tempering treatments on microstructure and impact toughness?
The results suggested that fresh untempered martensite has a significant effect on the impact toughness of the material. It was also noticed that in some cases double tempering may not produce a fully tempered microstructure and a third tempering may be useful.
Research question 4 has been studied in paper C:
4. How does holding time and temperature during tempering affect the transformation of retained austenite?
It was seen that the retained austenite can transform either by formation of martensite on cooling after holding at the tempering temperature or by transformation to ferrite and carbides during holding. Martensite formation on cooling was the main mechanism except for long time tempering at 600°C.
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CONCLUSIONS 6 Conclusions
The effects of austenitisation temperature, holding time, tempering temperatures and multiple tempering treatments were studied for the 5 wt.% Cr cold work tool steel ‘Caldie’. The following was concluded:
Effects of austenitisation temperature: Austenitisation at 1075°C caused significant grain coarsening compared to austenitisation at 1020°C or 1050°C. A higher austenitisation temperature led to a higher alloying content in the matrix and lower Ms and Bs temperatures. Thus, a higher volume fraction of retained austenite formed.
Tempering: Tempering at 200°C resulted in high contents of retained austenite, good impact toughness but low compressive yield strength. Tempering at 525°C resulted in lower amounts of retained austenite (<2 vol.%), sufficient hardness and high compressive strength. The presence of significant amounts of fresh martensite after a single tempering resulted in low impact toughness. A second tempering reduced the content of fresh martensite and increased the toughness. A triple tempering procedure can be needed when austenitising at high temperatures to decrease the level of fresh martensite in the tool microstructure to ensure sufficient toughness.
Transformation of retained austenite: Retained austenite transformation to martensite occurred on cooling to ambient temperature after tempering at 525°C. The carbon content of retained austenite decreased as a result of precipitation of carbides during holding at the tempering temperature. This resulted in destabilisation of the austenite and an increase of the martensite start temperature. Long-time tempering at 600°C resulted in transformation to ferrite and carbides during the tempering treatment.
Recommended heat treatment procedure: Austenitisation at 1020°C or 1050°C, followed by double tempering at 525°C produced a good combination of mechanical properties.
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FUTURE WORK 7 Future work
In this thesis, an attempt was made to understand the heat treatment process for a 5 wt.% Cr cold work tool steel. Some fundamental questions regarding austenitisation and tempering treatments and retained austenite transformation were answered. Naturally, some new questions were raised and interesting areas for further studies were revealed. The future work suggested below could contribute to further knowledge of importance both from the academic and industrial points of view.
It was not possible to conclusively identify the nature of blocky regions or directly observe carbide precipitation in retained austenite for short aging times by SEM. It would thus be useful to use techniques such as Transmission Electron Microscopy (TEM) and the SEM-Electron Back Scattered Detector (EBSD) for more detailed microstructural characterisation. The EBSD technique could also be used to measure the effective grain size of the martensitic microstructure which could then be related to the prior austenite grain size and effects on toughness.
The present study has illustrated that it is essential to minimize the content of retained austenite. It would therefore be of interest to investigate the effects and possibilities of applying cryogenic heat treatments and also effects of quenching rates on microstructure and mechanical properties.
Finally, the effects on industrially relevant properties of the possible alternative heat treatment parameters investigated during this study needs to be validated. An appropriate way would be to perform semi industrial punch testing of Caldie after different heat treatments. In this way, the work would gain industrial relevance and testing would be helpful in predicting tool behaviour during operation.
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REFERENCES 8 References
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SUMMARIES OF APPENDED PAPERS 9 Summaries of appended papers
9.1 Paper A
Effect of austenitisation temperature and multiple tempering on the microstructure and impact toughness of a 5 wt.% Cr cold work tool steel.
The effects on microstructure and properties of austenitisation at 1020°C, 1050°C or 1075°C followed by single, double or triple tempering treatments at 525°C were studied for a 5 wt.% Cr cold work tool steel. The transformation behaviour during quenching and tempering treatments was evaluated by dilatometry and correlated to the resulting microstructures. Impact toughness testing was performed after single, double and triple tempering treatments.
With higher austenitisation temperature, the martensite and bainite start temperatures were lowered resulting in larger volume fractions of retained austenite. The austenite formed martensite on cooling after tempering which was detrimental to toughness. Double or triple tempering treatments therefore resulted in higher un-notched impact toughness and a triple tempering was required to guarantee a fully tempered microstructure for the two highest austenitisation temperatures.
9.2 Paper B
Effect of austenitisation and tempering on the microstructure and mechanical properties of a 5 wt.% Cr cold work tool steel.
Effects of austenitisation at 1020°C, 1050°C or 1075°C followed by tempering at temperatures from 200°C to 600°C were investigated with an aim of understanding the influence on microstructure and mechanical properties. Microstructures were characterised and hardness, compressive yield strength and un-notched impact toughness were evaluated.
Austenitisation at 1075°C resulted in the highest amount of retained austenite, the highest hardness after tempering at 525°C and the lowest impact toughness. The lower impact toughness was due to coarsening of the martensitic microstructure. When the steel was tempered at 200°C, higher impact toughness and higher volume fractions of retained austenite were observed.
It was concluded that the best combination of mechanical properties is achieved by austenitisation at 1020°C or 1050°C followed by tempering at 525°C.
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SUMMARIES OF APPENDED PAPERS 9.3 Paper C
Retained austenite transformation during heat treatment of a 5 wt.% Cr cold work tool steel.
Retained austenite transformation was studied for a 5 wt.% Cr cold work tool steel tempered at 525°C and 600°C followed by cooling to room temperature. Transformation behaviour during holding at the tempering temperature and on cooling was evaluated. Microstructures were characterised and the lattice parameter of retained austenite was measured.
Tempering at 525°C showed retained austenite transformation to martensite on cooling. Longer holding times resulted in higher martensite start and finish temperatures while the lattice parameter of retained austenite decreased. These results suggested a depletion of carbon due to precipitation of carbides in retained austenite. Tempering at 600°C resulted in precipitation of carbides in retained austenite followed by transformation to ferrite and carbides.
It was concluded that there are two mechanisms of retained austenite transformation occurring depending on tempering temperature and time. The standard tempering treatment and the possibilities of achieving a fully tempered martensitic microstructure by alternative heat treatments were discussed based on these findings.
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Licentiate Thesis Licentiate Technology Production 18 No. 2017 Microstructure and mechanical and mechanical Microstructure cold of a 5 wt.% Cr properties of heat steel - Influence tool work treatment procedure. Muhammad Arbab Rehan
MUHAMMAD ARBAB REHAN MICROSTRUCTURE AND MECHANICAL PROPERTIES OF A 5 WT.% CR COLD WORK TOOL STEEL - INFLUENCE OF HEAT TREATMENT PROCEDURE. 2017 NO.18 ISBN 978-91-87531-56-9 (Printed) ISBN 978-91-87531-56-9 (Electronic) ISBN 978-91-87531-55-2